Ferritic stainles steel for exhaust gas path members

ABSTRACT

Provided is a stainless steel for exhaust gas path members, which has a composition including, in terms of % by mass, at most 0.03% of C, at most 1% of Si, at most 1.5% of Mn, at most 0.6% of Ni, from 10 to 20% of Cr, from more than 0.5 to 0.7% of Nb, from 0.05 to 0.3% of Ti, from more than 1 to 2% of Cu, at most 0.2% of V, at most 0.03% of N, from 0.0005 to 0.02% of B, and optionally at most 0.1% of Al, and further optionally at least one of Mo, W, Zr and Co in an amount of at most 4% in total, with a balance of Fe and inevitable impurities, and which has a texture where the Cu phase and the Nb compound phase having a major diameter of at least 0.5 μm are controlled to be in an amount of at most 10 grains/25 μm 2  each. The stainless steel exhibits excellent thermal fatigue resistance when applied to exhaust gas path members of both cases where the maximum ultimate temperature is high and low, and has excellent low-temperature toughness.

TECHNICAL FIELD

The present invention relates to a ferritic stainless steel for use as exhaust gas path members such as typically exhaust manifolds, catalyst converter casings (outer casings), front pipes and center pipes, and to an automobile exhaust gas path member using the same.

BACKGROUND ART

SUS444-type materials of good heat resistance are much used for exhaust gas path members such as exhaust manifolds, catalyst converter casings, front pipes, center pipes and the like. As a material improved in point of the high-temperature oxidation resistance and the high-temperature strength thereof within a high-temperature range over 700° C., Patent References 1 and 2 disclose a ferritic stainless steel with approximately from 1 to 2% by mass of Cu added thereto. Cu in the steel precipitates as a Cu phase when heated, and acts to enhance the high-temperature strength and the thermal fatigue resistance of the steel. The Cu-containing steel of the type is suitable to exhaust gas path members to be connected to engines of a type where the exhaust gas temperature is high.

Patent Reference 1: WO03/004714

Patent Reference 2: JP-A 2006-117985

Problems that the Invention is to Solve

Recently, the requirement for housing exhaust gas path members of automobile engines in a limited space owing to the increase in various devices to be mounted around engines is increasing, and a case of severely working and using them is increasing. Accordingly, even the members that are applied to engines where the exhaust gas temperature is not so high have become required to have extremely excellent thermal fatigue resistance and excellent low-temperature toughness.

For enhancing the high-temperature strength and the thermal fatigue resistance of ferritic stainless steel, there is known a method of adding a suitable amount of Cu to the steel, as in the above-mentioned Patent References 1 and 2; and in Patent Reference 2, in particular, there is employed a method of adding Nb to the steel in an amount of at most 0.6% by mass for the purpose of enhancing the high-temperature strength of the steel in a high-temperature range over 700° C. However, the present inventors' detailed studies have revealed that the Cu-containing steel in Patent References 1 and 2 may have good thermal fatigue resistance when the maximum ultimate temperature thereof is high (for example, from 200 to 900° C.), but when the maximum ultimate temperature thereof is low (for example, from 200 to 750° C.), its thermal fatigue resistance is somewhat inferior to that of SUS444-type materials. Accordingly, the steel of Patent References 1 and 2 is suitable to automobiles with high-power engines mounted therein where the exhaust gas temperature is high, but is not suitable so much for application to automobiles with small-sized engines mounted therein where the exhaust gas temperature is relatively low. Depending on the mode of its use, the exhaust gas temperature in a high-power engine may vary, and therefore for the exhaust gas path members, it is desired to use a material having good thermal fatigue resistance even in a case where the maximum ultimate temperature is low.

An object of the present invention is to provide a ferritic stainless steel capable of exhibiting excellent thermal fatigue resistance when applied to exhaust gas path members of both cases where the maximum ultimate temperature is high and low, and having excellent low-temperature toughness.

Means for Solving the Problems

As described in the above, in a case where the maximum ultimate temperature is high, for example, at 900° C. or more, the thermal fatigue resistance of steel can be enhanced through precipitation of a Cu phase therein. However, as a result of further studies, it has been clarified that the thermal fatigue resistance of steel in a case where the maximum ultimate temperature is low, for example, at around 750° C. or lower can be enhanced through control of the precipitation morphology of Nb in steel. Specifically, by controlling the precipitation morphology of the Cu phase and the Nb compound phase therein, a ferritic stainless steel can be realized capable of being well applicable to both cases where the maximum ultimate temperature is high and low.

The present invention provides a stainless steel for exhaust gas path members, which has a composition comprising, in terms of % by mass, at most 0.03% of C, at most 1% of Si, at most 1.5% of Mn, at most 0.6% of Ni, from 10 to 20% of Cr, from more than 0.5 to 0.7% of Nb, from 0.05 to 0.3% of Ti, from more than 1 to 2% of Cu, at most 0.2% of V, at most 0.03% of N, from 0.0005 to 0.02% of B, and optionally at most 0.1% of Al, and further optionally at least one of Mo, W, Zr and Co in an amount of at most 4% in total, with a balance of Fe and inevitable impurities, and having an [Nb] value, as defined by the following formula (2) or (3) in accordance with the [Ti] value thereof as defined by the following formula (1), falling within a range of from 0.5 to 0.65, and which has a texture where the Cu phase having a major diameter of at least 0.5 μm is controlled to be in an amount of at most 10 grains/25 μm² and the Nb compound phase having a major diameter of at least 0.5 μm is in an amount of at most 10 grains/25 μm²:

[Ti]=Ti−4(C+N)  (1),

when [Ti]≧0,[Nb]=Nb  (2),

when [Ti]<0,[Nb]=Nb+0.5[Ti]  (3).

The sites of Ti, C and N in the formula (1), and the site of Nb in the formula (2) and the formula (3) each are substituted with the value of the content of the corresponding element in terms of % by mass. Preferred embodiments of the exhaust gas path members include, for example, automobile exhaust manifolds, catalyst converters, front pipes and center pipes. Needless-to-say, the members may be applicable to any other various exhaust gas path members than those of automobiles.

The present invention has realized a ferritic stainless steel improved in point of both the thermal fatigue resistance thereof in a case where the maximum ultimate temperature is high (for example, 200 to 900° C.) and the thermal fatigue resistance thereof in a case where the maximum ultimate temperature is low (for example, 200 to 750° C.). Accordingly, the ferritic stainless steel of the invention is widely applicable to a case where the exhaust gas path members comprising it are used at a high exhaust gas temperature and to a case where they are used at a low exhaust gas temperature. In addition, the steel material satisfies the fundamental heat resistance (high-temperature oxidation resistance, high-temperature strength) required for automobile exhaust gas path members, and has excellent low-temperature toughness, and is therefore extremely useful for the recent exhaust gas path members that are required to be worked under severe conditions.

PREFERRED EMBODIMENTS OF THE INVENTION

The steel of the invention contains Cu and Nb, in which different types of precipitation phases of a Cu phase and an Nb compound phase are formed in actual service environments, and therefore, the steel exhibits excellent thermal fatigue resistance in both cases where the maximum ultimate temperature is high and low.

Various studies have revealed that, in the steel satisfying the composition mentioned below and having a texture state where the Cu phase having a major diameter of at least 0.5 μm is controlled to be in an amount of at most 10 grains/25 W 2 and the Nb compound phase having a major diameter of at least 0.5 μm is in an amount of at most 10 grains/25 μm², fine precipitates are formed sufficiently under heat in use of the steel, therefore bringing about remarkable enhancement of the thermal fatigue resistance of the steel. In other words, when both the precipitation phases of the Cu phase and the Nb compound phase having a major diameter of at least 0.5 μm previously exist in the steel in large quantities at a density of more than 10 grains/25 μm² therein, then fine precipitation phases could not form sufficiently under heat and the improvement in the stable thermal fatigue resistance of the steel could not be expected. In case where Cu or Nb is excessively in steel, overstepping the definition mentioned below, even when a coarse Cu phase or Nb compound phase exists in the material, there may be a possibility that the thermal fatigue resistance of the steel could be improved so far as fine precipitation phases could be formed therein. However, this case is unfavorable since the presence of the coarse crude precipitation phase therein may bring about a trouble of reducing the low-temperature toughness of the steel.

The Cu phase is a so-called ε-Cu precipitation phase, and this tends to readily grow in one direction and is generally in the form of rods. The Nb compound phase is a precipitation mainly comprising Fe₂Nb, and when containing Mo, it generally has a morphology of Fe₂(Mo,Nb). The Nb compound phase also tends to readily grow in one direction, and is generally in the form of rods. Accordingly, it is reasonable to evaluate the size of the precipitation phase as the major diameter thereof. Concretely, the major diameter of the precipitation appearing in the image projected through a transmission electronic microscope (TEM) (corresponding to the projection length in the image picture) may be taken as the major diameter of the phase. As to whether it is a Cu phase or an Nb compound phase, the projected phase may be identified with an analyzer (e.g., EDX) attached to TEM. Nb carbides and Nb nitrides are excluded from the Nb compound phase as referred to herein. The carbides and the nitrides are often in the form of masses or spheres, and from their forms, they may be relatively readily differentiated from the Fe₂Nb-type precipitation phase. In case where the differentiation is difficult from the forms, the phases may be identified with the above-mentioned analyzer (e.g., EDX).

In a case where the maximum ultimate temperature in use of the steel is around 900° C. or higher than it, Cu therein may sufficiently re-dissolve as a solid solution therein by that heating, therefore precipitating a fine Cu phase essentially at 500 to 700° C. Accordingly, the fatigue resistance in repeated heating (that is, the thermal fatigue resistance) of the steel can be thereby enhanced. On the other hand, in repeated heating where the maximum ultimate temperature is at most around 750° C. and is low, Cu could not sufficiently re-dissolve in the steel. Accordingly, the steel could not take a sufficient effect of enhancing the thermal fatigue resistance thereof owing to the fine precipitation of the Cu phase therein.

In the invention, the thermal fatigue resistance of the steel in the case where the maximum ultimate temperature is low and the Cu phase could not sufficiently improve the resistance by itself is compensated by the fine precipitation of the Nb compound phase in the steel. The Nb compound phase brings about precipitation reinforcement by heating at 700 to 750° C. though within an extremely short period of time. It has been found that the precipitation reinforcement phenomenon within the short period of time remarkably enhances the thermal fatigue resistance of the steel in a temperature range of from 200 to 750° C. At present, there are many unclear points of the mechanism, but it may be presumed that, as a result of the short-time precipitation reinforcement by the Nb compound phase, the steel could be free from balging to be caused by Ratchetting deformation or compression stress in the initial stage of repeated heating, and this could be advantageous for the thermal fatigue resistance of the steel where the maximum ultimate temperature is low.

The constitutive ingredients are described below.

C and N are generally said to be effective for enhancing the creep strength and other high-temperature strength of steel; however, when steel contains them excessively, then the oxidation resistance, the workability, the low-temperature toughness and the weldability of the steel worsen. In the invention, both C and N are limited to a content of at most 0.03% by mass each.

Si is effective for enhancing the high-temperature oxidation resistance of steel. In addition, it bonds to oxygen in the atmosphere in welding, therefore having the effect of preventing oxygen from penetrating into steel. However, when the Si content is excessive, then the hardness of the steel may increase and the workability and the low-temperature toughness thereof may worsen. In the invention, the Si content is limited to at most 1% by mass, and for example, it may be limited to a range of from 0.1 to 0.6% by mass.

Mn enhances the high-temperature oxidation resistance, especially the scale peeling resistance of steel. In addition, like Si, this bonds to oxygen in the atmosphere in welding, therefore having the effect of preventing oxygen from penetrating into steel. However, when added excessively, it worsens the workability and the weldability of steel. In addition, since Mn is an austenite-stabilizing element, a martensite phase may be readily formed when it is added too much, therefore causing a factor of reducing the workability of steel. Accordingly, the Mn content is limited to at most 1.5% by mass, more preferably at most 1.3% by mass. For example, the element may be defined to be from 0.1 to less than 1% by mass.

Ni is an austenite-stabilizing element, and when added excessively, it causes the formation of a martensite phase like Mn, therefore bringing about a factor of reducing the workability of steel. The Ni content is defined to be up to at most 0.6% by mass.

Cr stabilizes a ferrite phase and contributes toward enhancing the oxidation resistance of steel, an important property of high-temperature materials. However, too much Cr addition may make the steel material brittle and may cause the reduction in the workability of steel. Accordingly, the Cr content is to be from 10 to 20% by mass. Preferably, the Cr content is controlled in accordance with the temperature in use of the material. For example, in case where high-temperature oxidation resistance up to 950° C. is desired, the Cr content is preferably at least 16% by mass; and when the resistance up to 900° C. is desired, the Cr content may be from 12 to 16% by mass.

Nb is an element extremely effective for securing the high-temperature strength of steel in a high-temperature range over 700° C. It is believed that the Nb solid solution reinforcement in this composition system may greatly contribute toward the high-temperature strength enhancement. Nb fixes C and N and is therefore effective for preventing toughness reduction in steel. These effects of Nb are heretofore general ones; and further in the invention, the steel takes advantage of the fine precipitation of the Nb compound phase therein, therefore having another advantage of thermal fatigue resistance enhancement in a case where the maximum ultimate temperature is around at most 750° C. and is low (as described in the above). For sufficiently securing the effect of Nb, the Nb content must be more than 0.5% by mass, more effectively more than 0.6% by mass. However, too much addition of Nb worsens the workability and the low-temperature toughness of steel, and increases the high-temperature cracking sensitivity in welding; and therefore, the Nb content is limited up to at most 0.7% by mass.

On the other hand, Nb readily bonds to C and N. When Nb is consumed as its carbides and nitrides, then the high-temperature strength enhancement by the solid solution of Nb and the thermal fatigue resistance enhancement by the Nb compound phase may be insufficient. Accordingly, the [Nb] value of the composition, as defined by the following formula (2) or (3) in accordance with the [Ti] value thereof as defined by the following formula (1), or that is, the effective Nb amount in the composition is specifically defined.

[Ti]=Ti−4(C+N)  (1),

when [Ti]≧0, [Nb]=Nb  (2),

when [Ti]<0, [Nb]=Nb+0.5[Ti]  (3).

When the composition secures the Ti content over the amount capable of bonding to C and N, or that is, when the effective Ti amount [Ti] is not less than 0, then the Nb content value can be directly taken as the effective Nb amount [Nb] as in the formula (2). On the other hand, when the effective Ti amount [Ti] is less than 0, then the composition must secure an Nb content enough to compensate the effective Ti amount, and therefore, as in the formula (3), a value smaller than the Nb content is taken as the effective Nb amount [Nb].

In the invention, the Nb content is defined to fall within a range of from more than 0.5 to 0.7% by mass, and in addition, the effective Nb amount [Nb] is defined to fall within a range of from 0.5 to 0.65. In other words, the Nb content is severely defined within an extremely narrow range, and this is important to enhance the thermal fatigue resistance of the steel in a case where the maximum ultimate temperature is low, in addition to enhancing the high-temperature strength and the low-temperature toughness thereof.

Ti generally fixes C and N and is effective for enhancing the shapability of steel and preventing the toughness reduction in steel. Especially in the invention, the Ti content must also be severely controlled from the viewpoint of securing the effective Nb amount as so mentioned in the above. Concretely, the Ti content must be at least 0.05% by mass. However, addition of excessive Ti may worsen the surface property of steel owing to the formation of a large quantity of TiN, and may further have some negative influences on the weldability and the low-temperature toughness of steel. Accordingly, the Ti content is defined to be from 0.05 to 0.3% by mass.

Al is a deoxidizer, and is an element that acts for enhancing the high-temperature oxidation resistance of steel. In the invention, Al may be in the steel in a range of at most 0.1% by mass. Excessive Al, if any in steel, may form a large quantity of oxides in welding, and may act as a starting point of working cracks.

Cu is an important element for enhancing the high-temperature strength of steel. Specifically, as so mentioned in the above, the invention takes advantage of the fine dispersion precipitation phenomenon of the Cu phase of steel to thereby enhance the strength thereof at from 500 to 700° C. especially when the maximum ultimate temperature is not lower than around 900° C. and is high. Accordingly, the Cu content must be more than 1% by mass. However, too much Cu, if any in steel, may worsen the workability, the low-temperature toughness and the weldability of steel, and therefore, the Cu content is limited to be at most 2% by mass.

V contributes to enhancing the high-temperature strength of steel when added in combination with Nb and Cu. When existing along with Nb, V improves the workability, the low-temperature toughness, the resistance to grain boundary corrosion susceptibility, and the toughness of weld heat affected regions of steel. However, excessive addition of V rather impairs the workability and the low-temperature toughness of steel. Accordingly, the V content is limited to at most 0.2% by mass. Preferably, the V content is from 0.01 to 0.2% by mass, more preferably from 0.03 to 0.15% by mass.

B is effective in enhancing the resistance to secondary working brittleness of steel. The mechanism thereof is assumed to be due to the decrease in the solid solution C in grain boundaries or due to the strengthening of the grain boundaries. However, the addition of excessive B worsens the producibility and the weldability of steel. In the invention, the B content is from 0.0005 to 0.02% by mass.

Mo, W, Zr and Co are effective for enhancing the high-temperature strength of the ferritic stainless steel comprising the composition of the invention. If desired, at least one of these may be added to steel. However, too much addition may make the steel brittle, and therefore, when these elements are added, their total content is defined to be at most 4% by mass. More effectively, the total content is from 0.5 to 4% by mass.

The ferritic stainless steel having the above-mentioned composition can be produced according to an ordinary stainless steel smelting and casting process, and then, for example, this may be processed according to a cycle of “hot rolling→annealing→washing with acid” and optionally further according to one or more cycles of “cold rolling→annealing washing with acid”, thereby giving an annealed steel sheet having a thickness of, for example, from 1 to 2.5 mm or so. However, in the final annealing, it is important that, in the Nb precipitation temperature range and in the Cu precipitation temperature range, the steel is cooled at a suitable cooling speed. For example, as the final annealing condition employable herein, the steel material is heated at from 950 to 1100° C., preferably at from 1000 to 1100° C., and thereafter the mean cooling speed thereof from 1000 to 700° C. as a temperature range for Nb compound phase precipitation (when the heating temperature is lower than 1000° C., the mean cooling speed is from that heating temperature to 700° C.) is from more than 30 to 100° C./sec, and the mean cooling speed from 700 to 400° C. as a temperature range for Cu phase precipitation is from 5 to 50° C./sec. As a result of the above-mentioned composition control and the heat-treatment condition, a steel material (annealed steel sheet) can be obtained, having a texture condition where the Cu phase having a major diameter of at least 0.5 μm is controlled to be in an amount of at most 10 grains/25 μm² and the Nb compound phase having a major diameter of at least 0.5 μm is in an amount of at most 10 grains/25 μm². “Final annealing” as referred to herein is the final annealing to be carried out in a process of producing a steel material.

Using the annealed steel sheet, an exhaust gas path member is constructed. The tubular member is constructed as follows: The annealed steel sheet is roll-formed into a predetermined tubular form, then the butting portions of the two forms are welded to give a tube. For welding them, employable is any known tube-forming welding method of TIG welding, laser welding, high-frequency welding or the like. Thus obtained, the steel tube is, if desired, processed for heat treatment and washing with acid, and then worked into an exhaust gas path member.

EXAMPLES

A ferritic stainless steel having the composition shown in Table 1 was produced through smelting and casting, and then processed according to a process of “hot rolling→annealing/washing with acid→cold rolling→final annealing/washing with acid” to give an annealed steel sheet having a thickness of 2 mm. A part of the cast slab is hot-forged into a round rod having a diameter of about 25 mm, and this was processed for final annealing. The final annealing of the sheet and the final annealing of the rod were as follows, except for Steel No. 19. The sheet or the rod was soaked at 1050° C. for 1 minute, and then cooled from 1000° C. to 700° C. at a mean cooling speed of from more than 30 to 100° C./sec, and further from 700° C. to 400° C. at a mean cooling speed of from 5 to 50° C./sec. The final annealing of Steel No. 19 was as follows: From 1000° C. to 700° C., the mean cooling speed was controlled to be from 10 to 20° C./sec, and the other condition was the same as that of the other samples (common to both the sheet material and the rod material).

TABLE 1 Chemical Composition (mas. %) Group Steel C Si Mn Ni Cr Nb Ti Al Cu V N B Others [Ti] [Nb] Steels of the A 0.005 0.08 0.17 0.13 17.04 0.53 0.14 0.06 1.37 0.07 0.011 0.0018 Mo: 0.20 0.076 0.530 Invention B 0.007 0.12 0.25 0.16 17.00 0.55 0.16 0.03 1.40 0.06 0.010 0.014 — 0.092 0.550 C 0.008 0.15 0.31 0.12 17.05 0.61 0.12 0.02 1.51 0.05 0.009 0.0026 — 0.052 0.610 D 0.006 0.22 0.45 0.11 17.41 0.65 0.11 — 1.39 0.07 0.011 0.0019 — 0.042 0.650 E 0.005 0.89 0.23 0.09 11.49 0.63 0.06 0.02 1.41 0.05 0.007 0.0006 — 0.012 0.630 F 0.012 0.07 1.01 0.21 19.95 0.57 0.16 0.03 1.72 0.06 0.012 0.0015 — 0.064 0.570 G 0.021 0.06 0.14 0.58 16.88 0.62 0.28 0.02 1.42 0.07 0.022 0.0013 — 0.108 0.620 H 0.008 0.11 0.27 0.13 17.08 0.52 0.14 0.08 1.06 0.06 0.009 0.0052 Mo: 2.86 0.072 0.520 I 0.006 0.14 0.24 0.15 17.25 0.63 0.15 — 1.33 0.05 0.010 0.0016 Co: 3.09 0.086 0.630 J 0.005 0.15 0.09 0.11 17.61 0.55 0.16 0.02 1.22 0.07 0.009 0.0017 W: 2.48 0.104 0.550 K 0.006 0.16 0.22 0.14 17.21 0.61 0.14 0.03 1.26 0.06 0.010 0.0016 Zr: 0.59 0.076 0.610 L 0.009 0.23 0.32 0.16 16.88 0.51 0.15 0.04 1.21 0.03 0.009 0.0013 Mo: 1.51, 1.95 0.078 0.510 Comparative M 0.002 0.33 0.25 0.02 17.80 0.30 0.10 — 1.66 0.03 0.003 0.0005 — 0.080 0.300 Steels N 0.009 0.25 0.11 0.09 16.08 0.38 0.25 0.02 1.40 0.05 0.008 0.0010 — 0.182 0.380 O 0.028 0.33 0.31 0.11 17.40 0.45 0.07 — 1.48 0.04 0.021 0.0020 — −0.126 0.387 P 0.009 0.35 0.32 0.31 17.24 0.74 — 0.02 2.48 0.01 0.013 — — −0.088 0.696 Q 0.006 0.08 0.22 0.15 17.08 0.47 0.10 — 1.35 0.07 0.011 0.0010 Mo: 0.15 0.032 0.470 R 0.012 0.40 0.70 0.22 18.28 0.50 — — 0.24 0.01 0.021 0.0020 Mo: 1.94 −0.132 0.434 Underline: Outside the Scope of the Invention

The rolling direction of the sheet material and the longitudinal direction of the rod material are referred to as L direction. After the final annealing, the sheet material and the rod material were analyzed for the metal texture in the cross direction thereof cut perpendicularly to the L direction. Using a transmission electronic microscope (TEM), the Cu phase and the Nb compound phase were analyzed for their size, and the number of the Cu phase grains and the Nb compound phase grains having a major diameter of at least 0.5 μm seen per 25 μm² was counted. At least 10 viewing sites were analyzed in one sample, and the data were averaged. The type of the precipitation phase was identified by quantifying Fe, Nb, Mo and Cu with EDX (energy dispersion fluorescent X-ray analyzer) attached to TEM. In analyzing the fine precipitation phase, the constitutive elements of the steel base were detected together with the phase-constituting elements. Therefore, of the found data of the above-mentioned four elements to which the precipitation phase was targeted, the value of at least 50% by mass for Cu was identified as the Cu phase; and the value of at least 30% by mass of Nb was identified as the Nb compound phase. The samples where the Cu phase having a major diameter of at least 0.5 μm was in an amount of at most 10 grains/25 μm are O (good), and the others are x (not good), as in the column of the Cu phase in Table 2. The samples where the Nb compound phase having a major diameter of at least 0.5 μm was in an amount of at most 10 grains/25 μm² are O (good), and the others are x (not good), as in the column of the Nb compound phase in Table 2. There was found no difference in the test results between the sheet material and the rod material of every steel; and therefore, the precipitation phase evaluation shown in Table 2 applies to both the sheet material and the rod material of steels.

The sheet material was tested in an impact test for the low-temperature toughness thereof. A V-notched impact test piece was so collected that the direction in which the impact is to be given to the piece could be the rolling direction of the sheet. This was tested in an impact test of JIS Z2242 at a pitch of 25° C. in a range of from −75 to 50° C., and the ductility/brittleness transition temperature of the sample was determined. The samples of which the transition temperature was lower than −25° C. (the samples showing ductile fracture even at −25° C.) were evaluated as O (good), and the others were evaluated as x (not good).

The rod material was tested in a thermal fatigue test for the thermal fatigue resistance thereof at 200 to 750° C. and at 200 to 900° C. This was cut into a round rod test piece having a diameter of 10 mm and a parallel portion length of 20 mm between the gauge marks given thereto (the length between the gauge marks was 15 mm), and was notched to have R=5.7 mm and have a diameter of 7 mm at the center position between the gauge marks. This was tested and evaluated in air under the condition mentioned below. The number of the repeated cycles at which the stress of the sample lowered to 75% of the stress at the start of cracking thereof was defined as the thermal fatigue life.

[Thermal Fatigue Resistance at 200 to 750° C.]

A heat cycle of “200° C.×0.5 minutes soaking→heating up to 750° C. at a heating speed of about 3° C./sec→750° C.×2.0 minutes soaking→cooling to 200° C. at a cooling speed of about 3° C./sec” with a restraint ratio (ratio of imparted strain to thermal expansion) of 25% was repeated in every sample. The samples having a thermal fatigue life of at least 1800 cycles were evaluated as O (good); those having a thermal fatigue life of from 1500 cycles to less than 1800 cycles were evaluated as Δ (relatively not good); and those having a thermal fatigue life of less than 1500 cycles were evaluated as x (not good) The samples with the evaluation “O” passed the test.

[Thermal Fatigue Resistance at 200 to 900° C.]

A heat cycle of “200° C.×0.5 minutes soaking→heating up to 900° C. at a heating speed of about 3° C./sec→900° C.×0.5 minutes soaking→cooling to 200° C. at a cooling speed of about 3° C./sec” with a restraint ratio (ratio of imparted strain to thermal expansion) of 20% was repeated in every sample. The samples having a thermal fatigue life of at least 900 cycles were evaluated as O (good); and those having a thermal fatigue life of less than 900 cycles were evaluated as x (not good) The samples with the evaluation “O” passed the test.

The results are shown in Table 2.

TABLE 2 Precipitation Phase Low-Temperature Thermal Fatigue Resistance Group No. Steel Cu Phase Nb Compound Phase Toughness 200-750° C. 200-900° C. Samples of the 1 A ∘ ∘ ∘ ∘ ∘ Invention 2 B ∘ ∘ ∘ ∘ ∘ 3 C ∘ ∘ ∘ ∘ ∘ 4 D ∘ ∘ ∘ ∘ ∘ 5 E ∘ ∘ ∘ ∘ ∘ 6 F ∘ ∘ ∘ ∘ ∘ 7 G ∘ ∘ ∘ ∘ ∘ 8 H ∘ ∘ ∘ ∘ ∘ 9 I ∘ ∘ ∘ ∘ ∘ 10 J ∘ ∘ ∘ ∘ ∘ 11 K ∘ ∘ ∘ ∘ ∘ 12 L ∘ ∘ ∘ ∘ ∘ Comparative 13 M ∘ ∘ ∘ x ∘ Samples 14 N ∘ ∘ ∘ x ∘ 15 O ∘ ∘ ∘ Δ ∘ 16 P x x x ∘ ∘ 17 Q ∘ ∘ ∘ Δ ∘ 18 R ∘ ∘ ∘ Δ ∘ 19 B ∘ x x Δ ∘

Table 2 confirms that the samples of the invention, having the chemical composition as defined in the invention and satisfying the precipitation morphology of the Cu phase and the Nb compound phase, were improved in point of both the thermal fatigue resistance under the high maximum ultimate temperature condition (200 to 900° C.) and the thermal fatigue resistance under the low maximum ultimate temperature condition (200 to 750° C.), and had good low-temperature toughness.

As opposed to these, in the comparative samples, Nos. 13 to 15 and 17 in which the Nb content was low and the effective Nb amount [Nb] was also low, the formation of the fine Nb compound phase was unsatisfactory at the low maximum ultimate temperature of 750° C., and therefore the samples had poor thermal fatigue resistance at 200 to 750° C. No. 16 contained Cu and Nb excessively, but its thermal fatigue resistance was bettered though it contained many coarse Cu phase grains and Nb compound phase grains. However, its low-temperature toughness was poor. No. 18 is an ordinary steel corresponding to SUS444, in which the Cu content was low but the Mo content was high, and therefore, its thermal fatigue resistance at 200 to 900° C. was good. However, since the effective Nb amount in this was low, its thermal fatigue resistance at 200 to 900° C. was not bettered. No. 19 is a steel having the composition defined in the invention; however, in its final annealing, the cooling speed within the temperature range for Nb compound phase precipitation was too late, and coarse Nb compound phase grains were formed. As a result, in the following heating process, fine Nb compound phase grains could not be sufficiently precipitated in this, and therefore the thermal fatigue resistance at 200 to 750° C. of this sample was poor. In addition, owing to the influence of the coarse Nb compound phase grains thereon, the low-temperature toughness of the sample was also poor. 

1. A stainless steel for exhaust gas path members, which has a composition comprising, in terms of % by mass, at most 0.03% of C, at most 1% of Si, at most 1.5% of Mn, at most 0.6% of Ni, from 10 to 20% of Cr, from more than 0.5 to 0.7% of Nb, from 0.05 to 0.3% of Ti, from more than 1 to 2% of Cu, at most 0.2% of V, at most 0.03% of N, from 0.0005 to 0.02% of B, with a balance of Fe and inevitable impurities, and having an [Nb] value, as defined by the following formula (2) or (3) in accordance with the [Ti] value thereof as defined by the following formula (1), falling within a range of from 0.5 to 0.65, and which has a texture where the Cu phase having a major diameter of at least 0.5 μm is controlled to be in an amount of at most 10 grains/25 μm² and the Nb compound phase having a major diameter of at least 0.5 μm is in an amount of at most 10 grains/25 μm²: [Ti]=Ti−4(C+N)  (1), when [Ti]≧0, [Nb]=Nb  (2), when [Ti]<0, [Nb]=Nb+0.5[Ti]  (3).
 2. The stainless steel for exhaust gas path members as claimed in claim 1, of which the composition further contains at most 0.1% by mass of Al.
 3. The stainless steel for exhaust gas path members as claimed in claim 1, of which the composition further contains at least one of Mo, W, Zr and Co in an amount of at most 4% in total.
 4. An exhaust gas path member as claimed in claim 1, wherein the exhaust gas path member is any of automobile exhaust manifolds, catalyst converters, front pipes and center pipes.
 5. The stainless steel for exhaust gas path members as claimed in claim 2, of which the composition further contains at least one of Mo, W, Zr and Co in an amount of at most 4% in total.
 6. An exhaust gas path member as claimed in claim 2, wherein the exhaust gas path member is any of automobile exhaust manifolds, catalyst converters, front pipes and center pipes.
 7. An exhaust gas path member as claimed in claim 3, wherein the exhaust gas path member is any of automobile exhaust manifolds, catalyst converters, front pipes and center pipes. 